Nickel base superalloys and turbine components fabricated therefrom

ABSTRACT

A nickel base superalloy suitable for the production of a large, crack-free nickel-base superalloy gas turbine bucket suitable for use in a large land-based utility gas turbine engine, comprising, by weight percents:  
     Chromium 7.0 to 12.0  
     Carbon 0.06 to 0.10  
     Cobalt 5.0 to 15.0  
     Titanium 3.0 to 5.0  
     Aluminum 3.0 to 5.0  
     Tungsten 3.0 to 12.0  
     Molybdenum 1.0 to 5.0  
     Boron 0.0080 to 0.01  
     Rhenium 0 to 10.0  
     Tantalum 2.0 to 6.0  
     Columbium 0 to 2.0  
     Vanadium 0 to 3.0  
     Hafnium 0 to 2.0 and  
     remainder nickel and incidental impurities.

[0001] The present invention relates to directionally solidifiednickel-base superalloys alloys having improved heat treatcharacteristics, good high temperature longitudinal and transverse creepstrength properties, good hot corrosion resistance and resistance tooxidation. The invention also relates to the use of the alloys in thefabrication of turbine components, particularly large turbine bucketsand turbine blades for aircraft engines.

BACKGROUND OF THE INVENTION

[0002] It is known to employ nickel base superalloys in the fabricationof aircraft engine components. To be acceptable, such alloys mustexhibit good castability with no heat treat cracking, good hightemperature longitudinal and transverse creep strength properties andgood hot corrosion resistance.

[0003] One such nickel base superalloy employed as a turbine bladingmaterial in aircraft engines is single crystal (SC) Rene N4 alloy. Aform of SC Rene N4 is described in U.S. Pat. No. 5,154,884 as anickel-base superalloy composition comprising, by weight, 7-12% Cr, 1-5%Mo, 3-5% Ti, 3-5% Al, 5-15% Co, 3-12% W, up to 10% Re, 2-6% Ta, up to 2%Cb, up to 3% V, up to 2% Hf, the balance being essentially nickel andincidental impurities. U.S. Pat. No. 5,399,313 describes a modifiedversion of SC Rene N4 as comprising, by weight, 9.5-10.0 Cr, 7.0-8.0 Co,1.3-1.7 Mo, 5.75-6.25 W, 4.6-5.0 Ta, 3.4-3.6 Ti, 4.1-4.3 Al, 0.4-0.6 Cb,0.1-0.2 Hf, 0.05-0.07 C and 0.003-0.005 B, the balance being nickel andincidental impurities.

[0004] Typically, aircraft engine blades are small, on the order of afew inches long, and weigh a few ounces, or a few pounds at most. Powerturbine buckets, by contrast, are typically up to about 36 inches long,and weigh up to about 40 pounds. It has been found that use of singlecrystal alloys for such large parts is impractical. A need exists for asuperalloy for use in the fabrication of large turbine blades whichexhibits good castability with no heat treat cracking, good hightemperature longitudinal and transverse creep strength properties andgood hot corrosion resistance. The present invention seeks to satisfythat need.

SUMMARY OF THE INVENTION

[0005] The present invention is directed to an alloy and hightemperature heat treatment for buckets fabricated from nickel basesuperalloys that will allow the buckets to be used for extended periods,typically up to about 72,000 hours in a power turbine. It is has beenfound that such an extended turbine life can be achieved ifapproximately 60-80% solutioning of the gamma-prime precipitates in thealloy occurs. The gamma-prime precipitates provide the strengtheningphase for the alloy. Moreover, it has been discovered according to theinvention that adjusting the level of boron in the alloy of theinvention to within the range of about 70-130 ppm, generally about80-130 ppm, more usually about 80-100 ppm (about 0.0080-0.01 weight %),for example about 90 ppm (about 0.009 weight %), results in a reductionin the incidence of heat treat cracking in the cast buckets.

[0006] In a first aspect, there is provided a nickel base superalloysuitable for the production of a large, sound, crack-free nickel-basesuperalloy gas turbine bucket suitable for use in a large land-basedutility gas turbine engine, comprising or consisting essentially of, byweight percents:

[0007] Chromium 7.0 to 12.0

[0008] Cobalt 5.0 to 15.0

[0009] Carbon 0.06 to 0.10

[0010] Titanium 3.0 to 5.0

[0011] Aluminum 3.0 to 5.0

[0012] Tungsten 3.0 to 12.0

[0013] Molybdenum 1.0 to 5.0

[0014] Boron 0.0080 to 0.013

[0015] Rhenium 0 to 10.0

[0016] Tantalum 2.0 to 6.0

[0017] Columbium 0 to 2.0

[0018] Vanadium 0 to 3.0

[0019] Hafnium 0 to 2.0 and

[0020] Remainder nickel and incidental impurities.

[0021] A typical nickel base alloy of the invention comprises orconsists essentially of, in weight percent:

[0022] Chromium 9.50-10.00

[0023] Cobalt 7.00-8.00

[0024] Aluminum 4.10-4.30

[0025] Titanium 3.35-3.65

[0026] Tungsten 5.75-6.25

[0027] Molybdenum 1.30-1.70

[0028] Tantalum 4.60-5.00

[0029] Carbon 0.06-0.10

[0030] Zirconium 0.01 max (no min)

[0031] Boron 0.008-0.010 (also expressed as 80-100 parts per million(ppm))

[0032] Iron 0.20 max (no min)

[0033] Silicon 0.20 max (no min)

[0034] Manganese 0.01 max (no min)

[0035] Copper 0.10 max (no min)

[0036] Phosphorus 0.005 max (no min)

[0037] Sulfur 0.003 max (no min)

[0038] Columbium 0.40-0.60

[0039] Oxygen 0.002 max (no min)

[0040] Nitrogen 0.0015 max (no min)

[0041] Vanadium 0.10 max (no min)

[0042] Hafnium 0.10-0.20

[0043] Platinum 0.15 max (no min)

[0044] Rhenium 0.10 max (no min)

[0045] Rhenium+Tungsten 6.25 max (no min)

[0046] Magnesium 0.0035 max (no min)

[0047] Palladium 0.10 max (no min)

[0048] Nickel Remainder

[0049] In a further aspect, there is provided a method of making a castand heat treated article such as a large power turbine bucket of anickel-base superalloy of the invention, wherein the article is heatedin an argon atmosphere or in vacuum to develop 60-80 percent solutioningof gamma prime precipitate, followed by cooling to room temperature.Typically, the article is heated to a temperature of about 2260°F.-2300° F., but at least about 25° F. below the incipient meltingtemperature of the superalloy. The article may be cooled by a furnacecool at a cooling rate of about 35° F./minute to 2050° F., followed bygas fan cooling at nominally 100° F./minute to 1200° F., and then anycooling rate to room temperature.

[0050] In yet a further aspect, the invention provides an article, suchas a large turbine bucket, produced according to the method of theinvention.

[0051] In a further aspect, there is provided a gas turbine enginecontaining an article of the present invention.

[0052] The alloy of the invention exhibits several advantages. First, at90-130 ppm boron the alloy of the invention has better castability (forlarge turbine buckets) than SC Rene N4 at 30-50 ppm boron. Secondly, at90-130 ppm boron in DS form the alloy of the invention has an improvedyield over SC Rene N4 at 30-50 ppm boron. In regard to “yield”, SC ReneN4 implies one grain per part. SC Rene N4 is typically used to makesmall turbine blades. As small parts go, it is possible to have a true“single crystal.” However, for large components, it is difficult toactually produce a part with only one grain. Thus, “yield” for a SC partwould be near zero (i.e. it is not possible to fabricate any). Bychanging the composition of SC Rene N4 primarily by adding more boron,it is possible to make a multi-grained DS component. This multi-grainedDS component is designed to accommodate many grains across thecross-section of the part. Made in this manner, the “yield” increases to80-100%.

[0053] Thirdly, at 90-130 ppm boron, the alloy of the invention hasnominally equivalent mechanical properties (in the longitudinaldirection) to the SC Rene N4 at 30-50 ppm boron. Fourthly, at 90-130 ppmboron, the alloy of the invention has better transverse creep propertiesthan SC Rene N4 at 30-50 ppm. Fifthly, at 90 ppm boron, the alloy of theinvention has better resistance against heat treat cracking than eitherthe SC Rene N4 at 30-50 ppm boron or the 130 ppm boron DS alloy of theinvention. The alloy with 130 ppm boron has a lower melting point(approx. 2301° F.) than DS Rene N4 or DS Rene N4 with 90 ppm boron (m.p.approx. 2315° F.), or SC Rene N4 which has a melting point near 2334° F.(Melting points: DS Rene N4 with 130 ppm boron—2301° F.; DS Rene N4 with90 ppm boron—2315° F.; SC Rene N4 with 30-50 ppm boron—2334° F.).

BRIEF DESCRIPTION OF THE DRAWINGS

[0054] The invention will now be described in more detail with referenceto the accompanying drawings, in which:

[0055]FIG. 1 is a series of plots showing the effect of differentprocessing conditions on crack length in a MS7001 H turbine bucket; and

[0056]FIG. 2 is a regression plot showing creep strength as a functionof temperature;

[0057]FIG. 3 is a regression plot showing transverse creep strength (%)as a function of boron content (ppm);

[0058]FIG. 4 is plot showing creep elongation as a function of testtemperature;

[0059]FIG. 5 is a plot showing the effect of varying amounts of boron onincipient melting of SC or DS Rene D4;

[0060]FIG. 6 shows a third and fourth stage bucket fabricated from analloy of the invention; and

[0061]FIG. 7 is a gas turbine engine showing the location where bucketsof the invention are used.

DETAILED DESCRIPTION OF THE INVENTION

[0062] It has been found, according to the invention, that increasingthe boron from about 30-50 ppm in the SC Rene N4 specification to nogreater than 130 ppm boron, along with several changes in partconfiguration, including bucket shape, essentially eliminates castingcracks in large turbine buckets. The additional boron may create a“M₅B₃” phase where M is Ni or Ni₅B₃ eutectic phase in the grainboundaries and elsewhere within the alloy matrix (as determined by AugerSpectrometry and Microdiffraction analyses), and the melting propertiesof the alloy have been attributed to the presence of a “M₅B₃” boronphase. The presence of this eutectic phase lowers the incipient meltingpoint (the point at which the metal starts to melt) from 2334° F. to2301° F. (as determined by Differential Thermal Analysis (DTA)). Thus,after application of a 2320° F. heat treatment (normal for SC Rene N4),the DS alloys begin to melt at locations within the eutectic pools wherethe boron as Ni₅B₃ is concentrated. Many of these eutectic pools are inthe grain boundaries, and can be more segregated than those eutecticpools elsewhere within the grains. When the eutectic melting starts andthe bucket cools back down to room temperature, a linear imperfectiondefined as a crack may be created. These cracks, called heat treatcracks, may be several inches long but may not be visible to the unaidedeye. The heat treat cracks may be found by use of fluorescent penetrantinspection (FPI), a nondestructive inspection technique.

[0063] The inventors have carried out work to determine parameters withrespect to the boron content of the alloy. It has been found that boronat 30-50 ppm in the alloy of the invention is not particularly suitablefor castability of large buckets. At this level of boron, a 2320° F.heat treatment fully solutions the gamma-prime phase and providesoptimum longitudinal mechanical properties for long bucket life.However, at this low level of boron, the transverse creep properties areless than optimum for large buckets.

[0064] In contrast, boron at 130 ppm in the alloy has been found to besuitable for castability, but is not particularly suitable for a fullsolution heat treatment. The melting point of such an alloy is reducedto about 2301° F., and the highest heat treatment that may be reliablyapplied is 2280° F. if melting is to be avoided. Heat treatment at atemperature of 2280° F. provides only about 60-80% solutioning of thegamma-prime phase, but this is generally acceptable for a full-lifebucket. Thus, the gamma-prime phase in the 130 ppm boron material cannotbe fully solutioned because the alloy starts to melt before fullsolutioning can be achieved.

[0065] The transverse creep properties are acceptable with this higherlevel of boron of 130 ppm. However, at this level of boron, a 5% failurerate for heat treat cracking has been observed.

[0066] It has been found that a level of boron of about 80-100 ppm, i.e.about 90±10 ppm, is optimum for castability. In order to improve thelongitudinal creep properties for an improved margin for bucket life, anincrease in the percent gamma-prime solutioning over about 60-80% isdesired. This may be possible due to the increase in melting temperaturefor the intermediate (about 90 ppm) boron level. In addition, this 90ppm level of boron provides a greater margin against heat treatcracking, and increases the yield during the solution heat treatmentoperation.

[0067] Castability experiments have been performed using the proceduredescribed in U.S. Pat. No. 4,169,742 (herein incorporated by reference).A master “lean” heat of DSN4 was formed, where B and Zr were removed,but otherwise the remaining elements (except for C and Hf) were the sameas in SC Rene N4 as described above. A three-level, four-factor designedexperiment (DOE) was then carried out. Castability was examined usingthe aforementioned castability test with the grain boundarystrengthening elements (& Ti) at the following levels (Zr was not variedbut kept at the lowest level), as shown in the Table below: WeightPercent of Elements at the 3 levels Desired for DOE Experiment ElementLow Level Medium Level High Level Carbon 0.06 0.10 0.14 Hafnium 0.250.45 0.65 Boron 0.0075 (75 ppm) 0.01 (100 ppm) 0.015 (150 ppm) Titanium3.37 3.50 3.65

[0068] It has been determined that castability is improved if Hf and Tiare run at their highest levels, but this also depends upon the Bcontent. The differences between C and B could not be fully ascertainedbecause this was not a full factorial experiment (which would have been3×3×3×1×3 or 81 experiments), and due to the limited ranges of carbon(0.14%−0.06%=0.08%) and boron (0.015%−0.0075%=0.0075%) versus ranges forhafnium (0.65%−0.25%=0.45%) and titanium (3.65%−3.37%=0.28%).

[0069] Hafnium (Hf) is known to cause casting defects known as “bands”,which are transverse linear indications as determined during FPIexamination. It has been determined that 0.75% Hf causes bands in low orhigh boron DS Rene N4 (boron 30-50 ppm—or 80-130 ppm), whereas 0.25weight % Hf and 0.45 weight % Hf resulted in no bands. From thestandpoint of acceptable transverse creep ductility, the lower level ofHf in production buckets is not allowed to fall below 0.15 weight %.Thus, for DS Rene N4, Hf is generally maintained in the range of about0.15-0.45 weight %.

[0070] Experiments have been carried out using controlled amounts ofboron and hafnium added to a baseline N4 master heat to determine theireffect on castability, expressed as total inches of crack length. Themaster heat composition was, by weight, 0.04%C, 9.77% Cr, 7.49% Co,5.92% W, 1.51% Mo, 4.21% Al, 3.37% Ti, 0.45% Nb, 4.71% Ta, 0.16% Hf,0.00% B, less than 0.005% Zr, balance Ni. The results for thin wallcastings (about 60 mils thick) and thick wall castings (about 120 milsthick) are shown in the chart below. The least amount of cracking(expressed as inches of crack) is best. “Inches of Crack Length fromCastability Test”

Other heats made by doping master lean heat

[0071] The chart above shows that thin wall versus thick wall data arecomparable, and that best castability is observed for DS Rene N4 with 40ppm (0.004%) boron and no Hf, OR 130 ppm (0.013%) boron and 0.45% Hf,indicating there is a “saddle point” in the data. No Hf is notconsidered to be acceptable as it may decrease transverse creepductility. It has been found that castability of the 90 ppm boron alloywith 0.15% Hf is improved over the castability of 130 ppm boron materialwith 0.15% Hf. Higher Hf levels may create transverse “bands” or dross.Banding as noted earlier is a known casting flaw, and “dross” is anonmetallic inclusion caused by a chemical reaction between dissolvedoxygen in the metal and free hafnium in the metal which combine to forma stable oxide such as HfO₂ (hafnium oxide). In either case, lower Hf(typically 0.15-0.45 weight %) is desirable in creating defect-freecastings.

[0072] The method of the invention includes a ramp heat treatment up tothe solution heat treatment temperature plus the post-solution heattreatment cooling rate down to room temperature. Four factors areimportant to achieving reduced heat treatment cracking. Each has beeninvestigated at two levels, as discussed below.

[0073] HIP temperature (2175° F. or 2225° F.);

[0074] solution heat treat temperature (2270° F. or 2290° F.);

[0075] post-solution heat treatment temperature cooling rate (slowfurnace cool at about 35° F./minute, or fast gas fan cool at about 150°F./minute, both followed by gas fan cooling from a temperature of about2050° F.); and

[0076] solution heat treatment atmosphere (vacuum or argon gas).

[0077] HIP or “hot isostatic pressing” is a means by which internalporosity in the casting can be closed by the application of externalpressure. This is achieved in a HIP vessel. The porosity is closed bythe application of temperatures in the range of 2175° F.-2225° F. and15,000 psi for an alloy like SC or DS Rene N4.

[0078] A heat treat temperature of 2290° F. was chosen as the highesttemperature possible for the solution heat treatment. The temperature of2290° F. was reached using part of a RAMP4 cycle to 2290° F., which isset forth in the Table which follows: Typical RAMP4 Solution HeatTreatment Cycle to 2300 F. Hold Heating Ramp Rate Temp. Hold Time RatePurpose/Results 25 F./minute 1400 F.  10 mm. — Stabilize, and beginintroducing 800 microns of argon gas. Not used if already running in a100% argon atmosphere. 25 F./minute 2225 F.   8 hour Increase tohomogenize 25 F./hour 2250 F.   4 hours Increase to homogenize 30F./hour 2280 F.   2 hours Increase to homogenize 10 F./hour 2290 F.   2hours Increase to homogenize 10 F./hour 2300 F. 0.5 hours Cool to RTAchieve final gamma-prime solutioning

[0079] This heating cycle was chosen because there was no evidence ofmelting or heat treat cracking using a variety of bucket or ingot sizes.For the 2290° F. solution cycle, that part of the RAMP4 cycle above(including up to 2290° F./2 hours) was chosen. A temperature of 2290° F.was chosen because previous work by the inventor showed that at 2300°F., recrystallized grain (RX) defects could form in DS Rene N4, and toavoid the RX grains the temperature would have to be lowered. Since itis only possible to control the temperature to within 10° F., atemperature of 2290° F. was chosen as the highest practical heattreatment temperature.

[0080] The second solution heat treatment temperature was 2270° F. Thiswas based upon metallography photographs showing the percent ofgamma-prime solutioning, and was considered to be the lowest acceptabletemperature capable of providing a full-life bucket.

[0081] The results are set forth in FIG. 1. Heat treating at 2270°F.±10° F. was equivalent to heat treating in the range of 2260-2280° F.,and heat treating at 2290F±10° F. was equivalent to heat treating in therange of 2280-2300° F.

[0082] A reason that it is difficult to determine what causes heat treatcracking is because the buckets cannot be examined at the solution heattreatment temperature to see if they are cracked. It is necessary tocool the buckets down to room temperature for examination. In addition,the section size of the bucket has some effect on residual stress, whichfurther complicates the heat treat cracking issue.

[0083] The HIP temperature was probably not significant because it iswell below the incipient melting temperature. Furthermore, the HIP cycleis also a thermal cycle and therefore can provide some homogenization tothe DS Rene N4. In this case, the 2225° F. cycle would provide morehomogenization than the 2175° F. cycle. But based upon the experimentalanalysis, it was shown the amount of homogenization provided by eitherHIP cycle is inadequate to influence the heat treat cracking.

[0084] In addition to the previous HIP and solution heat treat cycles,the cooling rate was believed to have an effect on heat treat cracking.To investigate this, two cooling rates were employed. The first rate wasproduced from a gas fan cool in the range of 100-150° F./minute, whichis available on most vacuum furnaces. The second rate was selectedbecause it was used during development trials, specifically from Ramp 4heat treatment where gas fan cooling was not available—only naturalcooling was available (called furnace cooling). Furnace cooling isachieved by just turning off the furnace and letting it cool naturally.In this case, the range was measured to be 35-75° F./minute.

[0085] Finally, the furnace atmosphere was felt to be important. Twoatmospheres are commonly available. The first is a vacuum atmospherewith some argon backfill, in the range of 400-800 microns. The secondatmosphere that is commonly employed (and was used in RAMP 4 heat treat)was 100% argon (not a vacuum).

[0086] The furnace environment during the heat treat experiment wasdetermined to be a minor factor. Initially, it was thought a vacuum orpartial vacuum environment could cause heat treat cracking byvolatilizing the grain boundary elements. In this instance, during avacuum heat treatment, some elements with a low vapor pressure can beremoved from the alloy, possibly leaving void spaces such as along agrain boundary (which could be interpreted as a crack). However, neitheratmosphere (vacuum with partial pressure argon or 100% argon) had asignificant effect on the heat treat cracking of the DS Rene N4 buckets.

[0087]FIG. 1 shows that the cooling rate has the greatest influence onthe heat treat cracking, followed closely by the solution heat treatmenttemperature (the greater the slope, the larger the effect). The othertwo factors—HIP temperature and furnace atmosphere—are considered to beminor factors. Thus, the slower cooling rate and the lower solution heattreatment temperature afforded the best results (least amount of heattreat cracking).

[0088] When the alloy is DS Rene N4 alloy with 130 ppm boron, theoptimum heat treatment includes a HIP cycle at 15,000 psi for 4 hours inthe range of 2175-2225° F. followed by a solution heat treatmenttemperature in the range of 2270° F. to 2290° F., followed by a furnacecool of about 35° F./minute to about 2050° F. and gas fan cooling toless than 1200° F., to prevent heat treat cracking.

[0089] The solution temperature had the largest effect on heat treatcracking, and is generally 2280° F.±10° F. (i.e. 2270° F.-2290° F.),more usually 2280° F. This provides for a lower incidence of heat treatcracking while still achieving adequate gamma-prime precipitatesolutioning.

[0090] The cooling rate is generally in the range of 25-45° F./minute,for example 35° F./minute. The gas fan cooling may be initiated when thetemperature reaches approximately 2050° F.±50° F.

[0091] The furnace atmosphere may be 100% argon, or vacuum plus argonpartial pressure (400-800 microns). Vacuum plus argon partial pressure(400-800 microns) is generally employed. The use of this small amount ofargon helps reduce the vaporization (depletion) of chromium during theheat treat cycle.

[0092] From this 130 ppm boron group, 1 cracked bucket occurred out of19 total, or a 5.2% failure rate, due to heat treat cracking. Part ofthe reason for this is the small margin of error between the heat treattemperature (2280° F.) and the incipient melting point of this alloy(2301° F.). The temperature difference between heat treat temperatureand melting point is 2301−2280° F.=21° F. This small margin is less thanthe error of thermocouples, which would approach 1% of the actualtemperature, or at 2280° F. it would be 22.8° F. This means the actualheat treat temperature could exceed the true melting point of the alloy,without the furnace operator's knowledge. If that happened, it wouldcause incipient melting, which in the presence of residual stress maylead to heat treatment cracks. This is compared to a margin of 54° F.for the 40 ppm boron material between the heat treat temperature and thepotential for incipient melting and heat treat cracking (2334° F.−2280°F.=54° F.)

[0093] The margin for temperature error with a 2280° F. heat treatmentis shown in the Table below. Incipient Aim Heat Melting Point TreatMargin for Temp. DSN4/GTD444 (° F., on Temperature Error during HeatAlloys heating) (° F.) Treatment (° F.) DSN4 w/31 ppm 2346 2280 66 BoronDSN4 w/36 ppm 2344 2280 64 Boron DSN4 w/40 ppm 2334 2280 54 Boron DSN4w/90 ppm 2311 2280 31 Boron DSN4 w/130 ppm 2301 2280 21 Boron

[0094] The advantage in going to an intermediate level of boron, such asin the 80-100 ppm range, is in the margin between incipient melting(when the alloy starts to melt) at the 2280F heat treat temperature. Forexample, at 130 ppm B, there is only 21° F. between the incipientmelting point and the 2280° F. heat treatment. This is not an acceptablerange, because the error due to the thermocouple (TC) alone is 22.8F (1%of 2280F). But at 90 ppm B the range between incipient melting and theheat treat temperature has increased to 31° F. Therefore, afteraccounting for 22.8° F. of TC error, there is still 8.2° F. oftemperature margin (31° F.−22.8° F.) between the incipient melting pointand the 2280° F. heat treat temperature. While 8.2° F. of margin is nota lot, it is an equivalent margin when compared to other high-technologySC or DS alloys.

[0095] Buckets from 90 ppm boron heats were successfully heat treated at2280° F. with 0% failure rate due to heat treat cracking. For the 90 ppmboron material, the melting point was determined to be 2311° F. Thus,with a heat treat temperature of 2280° F. the temperature differencebetween the heat treat temperature and the melting point is 2311−2280°F.=31° F. The temperature difference between the heat treat temperatureand the incipient melting point is greater than the thermocouple error(1% of 2280° F. or 22.8° F.), so there is less opportunity forunknowingly heat treating the buckets above their incipient meltingpoint, causing heat treat cracking.

[0096] It has been found that the amount of boron influences theincipient melting point of the alloy, i.e. less boron is better. Theamount of boron additionally influences the transverse creep ductility,i.e. more boron is better (although boron does not influence thelongitudinal creep ductility). Moreover, a higher solution temperatureleads to more gamma prime solutioning, and more gamma prime solutioningleads to more longitudinal creep life. However, the solution temperatureinfluences the transverse creep ductility, whereby a lower temperatureis better.

[0097] Thus, optimization of the alloy requires transfer functions(equations) that describe these features in terms of controllablefactors. Additionally, creep strength and casting yield are not measuredin similar units. Therefore, the transfer function is expressed as apercentage of the best case for heat treat yield (100%) and creepstrength (100%). The transfer function generation is described below.

[0098] Heat treat yield is a function of two variables, boron contentand solution heat treatment temperature. If the B content is too high,incipient melting or heat treat cracking occurs at segregated areas inthe casting, resulting in scrap. If the solution heat treatmenttemperature is too high, incipient melting and recrystallization (RX)limit yield. Recrystallized grains result from a phase transformationwhere residual strains in the material on heating cause the formation ofstrain-free grains with little or no strength, i.e. critical defects.The following spreadsheet shows the data used to generate Heat TreatYield Transfer Function Equation 1: Heat Treat Yield Boron (B) (Percent)Temp. (F.) Content (ppm) 100 2280 40 50 2292 130 50 2310 40 90 2280 1300 2327 40 0 2310 130

[0099] Regression with the data leads to the following regressionequation:

Heat Teat Yield=5448−2.34(Temp)−(0.340)*(Boron content)  Eq. 1

[0100] This is the first transfer function for yield.

[0101] A statistical analysis was conducted for the data, resulting inthe following standard tables: Predictor Coef StDev T P VIF Constant5448.0 671.8 8.11 0.004 Temp −2.3353 0.2907 −8.03 0.004 1.1 B −0.33980.1117 −3.04 0.056 1.1

[0102] S=11.59 R-Sq.=95.6% R-Sq. (Adj)=92.6%

[0103] (R-Sq=R² or R squared; adj means Adjusted)

[0104] The next transfer function is for longitudinal creep strength.This is a function of gamma-prime precipitate solutioning versus thesolution heat treatment temperature, as the only way to get 100% creepstrength is to fully solution the material. The following is datarelating the percent of full creep strength versus the heat treattemperature for DS Rene N4: Creep Heat Treat Strength Temperature(Percent) (F.) 100 2320 90 2300 60 2280 40 2215

[0105] The longitudinal creep strength is in percent of maximumobtainable, and the heat treatment temperature (t) is the solution heattreatment temperature in degrees F.

[0106] The data was used to solve for Equation 2 (see the RegressionPlot in FIG. 2). The curve has the correct dependency of creep strengthon solution heat treatment temperature. It will be noted that as-cast DSRene N4 has about 40% of the possible creep strength and that solutionheat treatment of DS Rene N4 at 2320° F. gives 100% creep strength. Thisis the second transfer function.

[0107] A further important feature of the alloy is creep strengthtransverse (transverse creep strength) to the grain boundaries. This isimportant in the tip shroud and other areas where loading is not in aradial direction on he part. The following data was extracted fortransverse creep strength: Percent of Transverse Creep Boron ContentStrength. (ppm) 50 40 100 80 80 130 90 100

[0108] This information created a non-linear regression plot as shown inFIG. 3. Equation 3 is:

Y=−40.7431+2.9113X−1.54E−02X ²

[0109] The three transfer functions (equations) can be solvedsimultaneously using an optimization spreadsheet shown below: MULTIPLERESPONSE OPTIMIZATION

[0110] The solutions with respect to Heat Treat Yield, LongitudinalCreep and Transverse Creep Strength were: Needs Heat Treat Yield 1 1 2 23 3 Longitudinal Creep Strength 2 3 1 3 1 2 Transverse Creep Strength 32 3 1 2 1 Optimize B ppm 40 40 94.5 94.5 40 94.5 Temp F 2280 2280 22962280 2296 2280

[0111] A “1” means optimization on this need first, followed by “2” andfinally “3”.

[0112] This results in an optimized alloy with a boron content of94.5+/−10 ppm and a heat treatment temperature of 2280±20° F.

[0113]FIG. 4 is plot showing creep elongation as a function of testtemperature. FIG. 5 is a plot showing the effect of varying amounts ofboron on incipient melting of SC or DS Rene N4.

[0114]FIG. 6 shows a third and fourth stage bucket fabricated from analloy of the invention. FIG. 7 is a gas turbine engine showing thelocation where buckets of the invention are used.

[0115] While the invention has been described in connection with what ispresently considered to be the most practical and preferred embodiment,it is to be understood that the invention is not to be limited to thedisclosed embodiment, but on the contrary, is intended to cover variousmodifications and equivalent arrangements included within the spirit andscope of the appended claims.

What is claimed is:
 1. A nickel base superalloy suitable for the production of a large, crack-free nickel-base superalloy gas turbine bucket suitable for use in a large land-based utility gas turbine engine, comprising, by weight percents: Chromium 7.0 to 12.0 Carbon 0.06 to 0.10 Cobalt 5.0 to 15.0 Titanium 3.0 to 5.0 Aluminum 3.0 to 5.0 Tungsten 3.0 to 12.0 Molybdenum 1.0 to 5.0 Boron 0.0080 to 0.0130 Rhenium 0 to 10.0 Tantalum 2.0 to 6.0 Columbium 0 to 2.0 Vanadium 0 to 3.0 Hafnium 0 to 2.0 and remainder nickel and incidental impurities.
 2. The nickel base superalloy according to claim 1, wherein boron is present in an amount of about 0.008-0.010 weight percent.
 3. The nickel base superalloy according to claim 1, wherein boron is present in an amount of about 0.009 weight percent.
 4. The nickel base superalloy according to claim 1, wherein hafnium is present in an amount of about 0.015-0.45 weight percent.
 5. A nickel base superalloy suitable for the production of a large, crack-free nickel-base superalloy gas turbine bucket suitable for use in a large land-based utility gas turbine engine, comprising, by weight percents Chromium 9.50-10.00 Cobalt 7.00-8.00 Aluminum 4.10-4.30 Titanium 3.35-3.65 Tungsten 5.75-6.25 Molybdenum 1.30-1.70 Tantalum 4.60-5.00 Carbon 0.06-0.10 Zirconium 0.01 max Boron 0.008-0.010 Iron 0.20 max Silicon 0.20 max Manganese 0.01 max Copper 0.10 max Phosphorus 0.005 max Sulfur 0.003 max Columbium 0.40-0.60 Oxygen 0.002 max Nitrogen 0.0015 max Vanadium 0.10 max Hafnium 0.10-0.20 Platinum 0.15 max Rhenium 0.10 max Rhenium+Tungsten 6.25 max Magnesium 0.0035 max Palladium 0.10 max Nickel Remainder
 6. A method of making a cast and heat treated article of a nickel-base superalloy, comprising the steps of: (a) providing a superalloy of the composition of claim 1; (b) heating the superalloy to develop at least 60 percent solutioning of gamma prime precipitate; and (c) cooling to room temperature.
 7. The method according to claim 6 wherein the article is heated to a temperature of about 2260° F.-2300° F. but at least about 25° F. below the incipient melting temperature of the superalloy.
 8. The method according to claim 6 wherein the article is cooled by a furnace cool at a rate of about 35° F./minute to about 2050° F.
 9. The method of claim 6 wherein the wherein the article is cooled by a gas fan cool at a rate of about 100-150° F./minute from below about 2050° F.
 10. The method of claim 6, wherein said heating is carried out in an argon atmosphere.
 11. The method of claim 6 wherein said heat treating comprises the steps of: (a) heating said article to a temperature of about 1400° F. at a rate of 25° F./minute and holding for about 10 minutes; (b) heating said article in (a) to a temperature of about 2225° F. at a rate of 25° F./minute and holding for about 8 hours; (c) heating said article in (b) to a temperature of about 2250° F. at a rate of 25° F./minute and holding for about 4 hours; (d) heating said article in (c) to a temperature of about 2280° F. at a rate of 30° F./minute and holding for about 2 hours; and (e) cooling to room temperature.
 12. The method according to claim 11 wherein the article is cooled by a furnace cool at a rate of about 35° F./minute to about 2050° F.
 13. The method of claim 11 wherein the wherein the article is cooled by a gas fan cool at a rate of about 100-150° F./minute from below about 2050° F.
 14. The method of claim 6 wherein said article is a large turbine bucket.
 15. The method of claim 6 wherein said article is a large aero engine turbine blade.
 16. An article produced by the method of claim
 6. 17. An article of claim 16 which is directionally solidified.
 18. An article of claim 16 which is conventionally cast.
 19. An article of claim 16 which is single crystal cast.
 20. A gas turbine engine containing an article of claim
 16. 